Solution to the Hole-Doping Problem and Tunable Quantum Hall

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Solution to the hole-doping problem and tunable quantum Hall effect in BiSe thin films 2

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JISOO MOON, Nikesh Koirala, Maryam Salehi, Wenhan Zhang, Weida Wu, and Seongshik Oh Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.7b04033 • Publication Date (Web): 09 Jan 2018 Downloaded from http://pubs.acs.org on January 11, 2018

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Solution to the hole-doping problem and tunable quantum Hall effect in Bi2Se3 thin films Jisoo Moon1, Nikesh Koirala1,†, Maryam Salehi2, Wenhan Zhang1, Weida Wu1, Seongshik Oh1*

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Department of Physics & Astronomy, Rutgers, The State University of New Jersey, Piscataway,

New Jersey 08854, U.S.A. 2

Department of Materials Science & Engineering, Rutgers, The State University of New Jersey,

Piscataway, New Jersey 08854, U.S.A.

*Correspondence should be addressed to [email protected] and +1 (848) 445-8754 (S.O.)

ACS Paragon Plus Environment

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ABSTRACT. Bi2Se3, one of the most widely studied topological insulators (TIs), is naturally electron-doped due to n-type native defects. However, many years of efforts to achieve p-type Bi2Se3 thin films have failed so far. Here, we provide a solution to this long-standing problem, showing that the main culprit has been the high density of interfacial defects. By suppressing these defects through an interfacial engineering scheme, we have successfully implemented ptype Bi2Se3 thin films down to the thinnest topological regime. On this platform, we present the first tunable quantum Hall effect (QHE) study in Bi2Se3 thin films, and reveal not only significantly asymmetric QHE signatures across the Dirac point but also the presence of competing anomalous states near the zeroth Landau level. The availability of doping tunable Bi2Se3 thin films will now make it possible to implement various topological quantum devices, previously inaccessible.

KEYWORDS. Topological insulator, Bi2Se3, Doping, Interface, Quantum Hall effect


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Hole (p) doping has been challenging in Bi2Se31–3, one of the most widely studied topological insulators (TIs)4–10. Unlike conventional semiconductor materials, the problem is complicated due to the presence of both surface and bulk states in topological insulators: we have to consider the doping problem of the surface and the bulk states separately. Both the surface and bulk states of Bi2Se3 have a strong tendency toward n-type due to its native n-type defects such as selenium vacancies1–3,11–17. In bulk crystals compensation dopants such as Ca and Mn can be used to convert the dominant carrier type from n- to p-type18–25. However, such a compensation doping scheme has not been successful in thin films of Bi2Se3. More specifically, we have tried various potential p-type dopants such as Zn, Mg, Ca, Sr, and Ba as compensation dopants, but none of them have so far led to p-type Bi2Se3 thin films; no p-type Bi2Se3 thin films have been demonstrated in the literature either. Only when they were quite thick (~200 nm), we were able to achieve p-type Bi2Se3 films through a complex process8. This difficulty in achieving p-type Bi2Se3 thin films has been puzzling, considering the very existence of p-type Bi2Se3 bulk crystals19,20. It may be suspected that this discrepancy in doping efficiency between thin films and bulk crystals could be due to the different growth conditions of the two systems, such as growth temperatures, considering that films are, in general, grown at much lower temperatures than bulk crystals. However, it should be noted that even p-type Bi2Se3 bulk crystals tend to become n-type when the crystals are made into thin flakes through, say, the Scotch tape method26. Bi2Se3 flakes can remain p-type only if they are relatively thick (> ~150 nm)24. All these observations provide evidence that whether they are thin films or bulk crystals, thickness critically affects the doping efficiency of the Bi2Se3 system, but the origin behind this problem has been a mystery, not to mention that a solution does not exist. Finding a solution to this pdoping mystery in Bi2Se3 thin films would be a major breakthrough toward the age of


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topological electronics, just as the solution to the long-standing p-doping problem in GaN has led to the solid-state-lighting revolution27. In order to better understand the hole-doping problem, we grew a series of Bi2Se3 thin films with different compensation doping schemes and different thicknesses and studied their carrier types and densities as shown in Figure 1. As for the compensation dopant, we have chosen Ca, but according to our comparison studies (Supporting Information Figure S3 and Section IV), other alkaline earth elements such as Mg and Sr might work as well. Figure 1a shows that when Bi2Se3 thin films are directly grown on Al2O3 (0001) substrates without any capping layer or a buffer layer, the films remain n-type regardless of the compensation-dopinglevel for film thickness up to 50 QL (quintuple layer; 1 QL is approximately 1 nm thick): as for how to control and estimate the Ca-doping level, see Supporting Information Section II. In order to check the effect of the air exposure, we subsequently grew another set of samples with Se capping layer, which was previously shown to be effective in protecting the surface of Bi2Se3 films from air exposure28. As depicted in Figure 1b, the Se capping does make a difference and 50 QL Bi2Se3 film on Al2O3 is now converted to p-type. This suggests that part of the reason why counter-doping for Bi2Se3 thin films becomes ineffective is due to air exposure, which could neutralize the easily-oxidizing Ca dopant as well as add extra n-type dopant such as water vapor10. However, even with the Se capping, the counter-doping still fails for much thinner (6 QL) films. This shows that air exposure is not the only reason for the failure of charge compensation doping with Ca, suggesting that the interface with the substrate is another potential source of the problem.


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As reported recently6, most of the native defects responsible for the high level of n-type carrier densities in Bi2Se3 thin films originate from the interface with commercially available substrates due to chemical and structural mismatch. This leads to high density of n-type carriers and prevents Bi2Se3 thin films from being hole-doped29. Therefore, we employ In2Se3 as a buffer layer that works as an excellent substrate for Bi2Se3 thin films, sharing chemical similarity and the same crystal structure. However, In2Se3 has multiple phases unlike Bi2Se3, and grows on Al2O3 in a disordered form. In order to overcome this problem, we grow a seed layer of 3 QL Bi2Se3 on Al2O3 substrate at 135 ℃ before growing In2Se3: for all subsequent layers, our typical growth rate is ~1 QL/min. The seed layer serves as a substrate for In2Se3, and it is followed by 20 QL of In2Se3 deposition at 300 ℃. Then, we anneal it to 600 ℃, which forces the Bi2Se3 layer to be diffused into In2Se3 layer and eventually vaporized out. As the last step in the buffer layer growth, 20 QL (Bi0.5In0.5)2Se3 (in short BIS) is deposited on the highly crystalline In2Se3 layer at 275 ℃. This 50% BIS layer provides Bi2Se3 not only with less lattice mismatch, ~1.6 %, than In2Se3 layer (~3.3%) but also with suppression of In diffusion into the (Ca-doped) Bi2Se3 film on top. Now, we grow Ca-doped Bi2Se3 film on the BIS at 275 ℃, then an atomically sharp and defect-free interface between Bi2Se3 and the BIS can be obtained. Utilizing the properlydesigned BIS buffer layer, we were able to eliminate most of these interfacial defects. The surface Fermi levels of these films were so low that they revealed a series of - previously inaccessible - topological quantum effects such as quantum Hall effect (QHE)6, quantized Faraday/Kerr rotation30, and quantum size effect during topological phase transitions31. The significantly reduced defect density in these films can be seen in the Hall resistance data of Figure 1c: an 8 QL Bi2Se3 thin film directly grown on Al2O3 (0001) shows an n-type sheet


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carrier density of 2.5 × 1013 /cm2, whereas a similar film on the BIS buffer layer exhibits only 1.0 × 1012 /cm2, 1/25 of the former value. Now, by employing both the BIS buffer layer and the Se capping layer, we were able to implement p-type Bi2Se3 thin films down to the thinnest topological regime (6 QL) as shown in Figure 1d. As summarized in Figure 1e, previous failures of counter-doping for thin Bi2Se3 films originate from two factors: first, neutralization of the counter-dopants due to air exposure on the top surface and second, high density of defects at the interface with the substrate. When Bi2Se3 films are grown directly on commercial substrates, there exists a high density of defects confined to the interface between the film and the substrate. These interfacial defects lead to high level of n-type carriers with the surface Fermi level located far above the bottom of the conduction band minimum2,8,32. Our current study suggests that this interfacial defect density on the commercial substrate is higher than the solubility limit of the compensation dopant so that the film remains ntype up to the maximum counter-doping (Figure 1f). On the other hand, with the BIS buffer layer, the interfacial defect density drops below the solubility limit of the compensation dopant and the Fermi level can be tuned below the Dirac point, leading to p-type Bi2Se3 films. We now provide more detailed Hall effect data for a range of Ca-doped Bi2Se3 films grown on the BIS buffer and the Se capping layers in Figure 2a. As the Ca doping level increases and compensates for the intrinsic n-type carriers, the (n-type) sheet carrier density gradually decreases, which appears as increasing negative slope of the Rxy(B) curve as illustrated in Figure 2c. i-ii. As we add more Ca to the films, the slope changes from negative (n-type) to positive (ptype), passing through a nonlinear n-p mixed regime: Figure 2c. iii-iv. 50 QL films enter the n-p mixed regime at 0.09% of Ca and become fully p-type at 0.13%. However, as the films get thinner, gradually higher level of Ca doping is required to reach the n-p mixed and the p regimes.


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The n-p mixed regime started at 0.13% for 20 QL, 0.3% for 10 QL, and 0.9% for 6 QL, and the p regime started at 0.3% for 20 QL, 2.0% for 10 QL and 6 QL. As summarized in the phase diagram in Figure 2b, these observations show that counter-doping becomes gradually more difficult, requiring a higher level of Ca doping, as Bi2Se3 films become thinner. This suggests that even in these interface-engineered Bi2Se3 films, the majority of defects still originate from the interfaces rather than from the bulk of the films. Even though p-type Bi2Se3 films are achieved for all thicknesses from 50 QL down to 6 QL at 2% of Ca, the films again become n-type with further increase (above 4.4%) in Ca doping except for 6 QL, which becomes insulating probably due to a disorder-driven topological phase transition31,33. This is because Ca doping inevitably introduces disorder and thus degrades both electronic and structural properties of the Bi2Se3 films. This can be seen in the carrier density and mobility vs Ca doping plots of Figure 3a and b. In Figure 3a, we plot the sheet carrier density (n2D) vs Ca doping for the 20 QL Bi2Se3 films grown on the BIS buffer with Se capping layers. For simplicity, we used the Hall resistance data at B = 0 T for the carrier density, and excluded the nonlinear n-p mixed regime because of ill-defined n2D in nonlinear Rxy(B) curves: more elaborate discussion of the data including the n-p mixed regime is given in Supporting Information Section III. It shows that the carrier type changes from n to p, and finally to n again at very high Ca concentration. In Figure 3b, we plot the corresponding mobility calculated from µ = 1 / (n2D · e · Rxx), where Rxx is the sheet resistance. The mobility shows a sharp drop as soon as the film becomes p-type, and gradually degrades with increasing Ca concentration in the pregime: as we further discuss with Figure 4d below, the sharp drop of the mobility through the charge neutral point seems to be an intrinsic property of the Bi2Se3 band structure, whereas its gradual decrease at higher Ca concentration is more due to disorder.


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This observation suggests that once Ca concentration goes beyond a certain solubility limit, Ca in Bi2Se3 stops acting as a p-type dopant and instead starts to act as n-type. This is well understood from the very fact that most defects in Bi2Se3 behave as n-type2. High level of Ca doping degrades not only the mobility but also the morphology, as can be seen in the Scanning Tunneling Microscope (STM) images of Figure 3c and d. The STM images show that the morphology of Bi2Se3 films degrades noticeably at high Ca concentration. The re-entrant n-type behavior beyond a solubility limit of Ca doping explains why Bi2Se3 thin films without the buffer layer cannot be turned into p-type via compensation doping: high density of (n-type) interfacial defects requires proportionately high density of compensation dopant to reach p-type, but beyond a certain doping density, the compensation dopant starts to act as n- instead of p-type, making it impossible to convert the film into p-type. Finally, these p-type Bi2Se3 thin films have allowed us to take the first doping-dependent QHE study in TI materials. Although gate-dependent QHE have been observed in (Bi1−xSbx)2Te3 thin films34 and in BiSbTeSe2 single crystal flakes35, and non-gated QHE in Bi2Se3 thin films6, doping-dependent QHE has not been available in any TI materials. Getting access to quantum regime in both n and p-doped TI films is critical for many topological quantum devices. For this study, we grew a series of 8 QL Bi2Se3 thin films with different Ca doping levels and measured their Hall effect up to a 34.5 T magnetic field with hand-patterned Hall bar geometry. In addition, we added MoO3 as the main capping layer under the Se capping as shown in Figure 4d, because MoO3 helps reduce the Ca doping level to reach the p-type6. Figure 4a shows Cadoping-dependent QHE in the Bi2Se3 films. The Hall resistance at high magnetic field is perfectly quantized at Rxy = h/e2 = 25.8 kΩ with vanishing sheet resistance on the n-side up to 0.08% of Ca, indicating the emergence of a well-defined chiral edge channel at high magnetic


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fields. Here, it is notable that only the 1st Landau level (LL) plateau (Rxy = h/(νe2) with ν = 1) is observable. The absence of higher level plateaus could be due to multiple factors such as the formation of non-chiral edge channels on the side surfaces36, the presence of electronic puddles11,37 and non-zero coupling between the transport channels of the top and bottom surface states4. As Ca doping is increased toward the Dirac point (or, equivalently, zeroth LL), while the Hall resistance gradually deviates from the quantized value, the sheet magnetoresistance changes drastically even with minute change in doping. At 0.09%, the Hall resistance at high magnetic field deviates just slightly from the quantized value, but the sheet magnetoresistance soars from close-to-zero to a large insulating value (~40 kΩ) at high magnetic fields as the doping changes from 0.08 to 0.09% even if the resistance remains nearly the same at zero magnetic field. This nominally implies that the dissipationless edge state is breaking down while the rest of the sample is insulating under high magnetic field. However, as we get even closer to the zeroth LL at 0.11%, the Hall resistance starts to decrease over the entire magnetic field but the sheet magnetoresistance plummets: this nominally implies that the - otherwise insulating - sample is now becoming conducting. The overall decrease in Hall resistance at 0.11% is likely due to the n-p mixing as we discussed above in Figure 2. On the other hand, the presence of both insulating and conducting - in terms of sheet magnetoresistance - states right next to each other near the zeroth LL is puzzling, and it suggests that the zeroth LL harbors multiple unidentified states, which would require in-depth future studies, both experimentally and theoretically. With further increase of Ca doping to 0.12 and 0.14%, the majority carrier type now converts to p-type. There are a couple of features noticeable in the p-type films. First, the p-type carrier density grows quickly to ~1012 /cm2 (Figure 4b) with a noticeable signature of n-p mixing:


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pure p-regime can be achieved with much higher (~1%) level of Ca doping, but only with about an order of magnitude higher carrier density than the pure n-regime (see Supporting Information Section V). Second, the signature of QHE is significantly degraded in the p-regime with the 1st LL barely observable only for 0.12%. Third, as can be seen in Figure 4c, the mobility sharply decreases as soon as the majority carrier type changes from n to p-type: this is consistent with what we observed in Figure 3. Although Ca-doping can degrade the mobility to some extent, it cannot be the main factor at these low doping levels: note that the mobility in Figure 4c remains almost constant or even grows a little with Ca-doping in the n-regime until it reaches the charge neutral point, after which the mobility sharply decreases. This suggests that the mobility drop right on the p-side must be of intrinsic origin. Both the sharp drop in mobility and the sharp increase in carrier density on the p-side are very likely due to the band structure effect. According to the band diagram of Bi2Se338, the Dirac point is very close to the bulk valence band and the surface band broadens noticeably on the p-side: both of these effects can account for the observed degradation of transport properties on the p-side of Bi2Se3 thin films. The near-absence of QHE and low mobility on the p-side of Bi2Se3 system is also consistent with the absence of Landau levels on the p-side as measured by a previous Scanning Tunneling Spectroscopy (STS) study39. In conclusion, we have provided a solution to the long-standing hole doping problem and presented the first doping-dependent QHE study in Bi2Se3 thin films. We have shown that the previous failures of hole doping efforts in Bi2Se3 films are due to: first, large density of interfacial defects from the substrates, and second, neutralization of the dopant from the air. By suppressing the interfacial defects via the (Bi0.5In0.5)2Se3 buffer layer and protecting the top surface with a capping layer, we were able to reach the hole regime of Bi2Se3 films down to the


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thinnest topological regime (6 QL) through compensation doping. In addition, in the doping dependent QHE study of 8 QL Bi2Se3 thin films we showed that the first (and only) Landau level is clearly observed on the electron side but quickly dissolves away on the hole side. Interestingly, at close proximity to the zeroth Landau level, we observed both the disintegration of the edge channel and the emergence of a conducting state, whose origins are unknown at the moment. Altogether, just as the solution to the p-doping problem of GaN has led to the LED lighting revolution, we believe that this solution to the p-doping problem of Bi2Se3 thin films will help usher in the age of topological quantum electronics.


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MBE growth Films were grown on 10 mm × 10 mm Al2O3 (0001) substrates (MTI Corporation) using a custom-built MBE system (SVTA) with base pressure less than 5 × 10-10 Torr. Substrates were cleaned ex situ by UV-generated ozone and in situ by heating to 750 ºC under oxygen pressure of 1 × 10-6 Torr. High purity (99.999%) Bi, In, Se, and Ca were thermally evaporated using Knudsen effusion cells for the film growth. Flux ratio of Se to Bi and In was maintained at above 10:1 to prevent the films from defects arising from Se deficiency. Source fluxes were calibrated in situ by quartz crystal micro-balance (QCM). For growth of In2Se3-(Bi0.5In0.5)2Se3 buffer layer and control of Ca flux relative to Bi (doping level), see Supporting Information Section I and II. MoO3 (for 8QL films) and Se were deposited at room temperature for capping on top of (CaxBi1x)2Se3

layers.

Transport measurements DC magneto-resistance and the Hall resistance measurements were carried out in a liquid He cryogenic system (AMI) with base temperature of 1.6 K. Perpendicular magnetic field (B) was applied to the films up to ± 9 Tesla using a solenoid superconductor coil. Electrical contacts were made by pressing thin In leads in van der Pauw geometry. Electrical current was generated and longitudinal, transverse voltages were measured for Rxx and Rxy using KE2400 sourcemeter and KE7001 switch matrix system. Raw data of Rxx(B) and Rxy(B) were symmetrized and antisymmetrized, respectively. STM measurements STM measurements were carried out at 77 K in an Omicron LT-STM with a base pressure of 8 × 10-12 Torr. Electrochemically etched tungsten tip was characterized on single


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crystal Au (111) surface. The samples, which were capped by Se in the MBE chamber, were exposed to air for several minutes during transfer to the STM chamber and the Se capping was decapped in a prep chamber, attached to the STM main chamber, with a base pressure of 2 × 1011

Torr. The samples were first sputtered with 500 eV argon ions for 15 minutes. Subsequently,

they were annealed up to ~200 °C to desorb the amorphous Se, after which they were immediately transferred into the STM head for measurements40.


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Figure 1. Ca doping study in Bi2Se3 thin films with different interfacial layers (at 1.5 K). (a-d) Hall resistance data, (e) summary table, and (f) schematic of energy band diagram with and without the interface layers. (a) Grown directly on Al2O3 (0001) substrates: the films remain ntype regardless of Ca-doping levels. (b) Grown on Al2O3 substrates with a Se capping layer: thick (50 QL) film can become p-type whereas thin (6 QL) film still remains n-type. (c) Effect of the BIS buffer layer on the interfacial defect density: the Bi2Se3 film grown on the buffer layer shows much lower carrier density than the one directly grown on Al2O3. (d) Achievement of ptype Bi2Se3 films down to 6 QL with the BIS buffer and the Se capping layers. (e) Summary


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table for the counter-doping effectiveness for different thicknesses and interfacial layers. (f) Schematic drawings showing the Fermi level tunability with Ca doping in Bi2Se3 films grown with and without the interfacial layers.

Figure 2. Hall resistance study for n- and p-type Bi2Se3 thin films (at 1.5 K). (a) Rxy(B) data vs Ca concentration and thickness for Ca-doped Bi2Se3 films with the BIS buffer and the Se capping layers. n-type (negative slope) and p-type (positive slope) curves are colored as red and blue, respectively, and the nonlinear, n-p mixed curves are colored as pink. As Ca concentration increases, all the films transition from n- to p-type through n-p mixed regime, and then eventually become n-type again except for the 6 QL film, which becomes insulating, instead. (b) A phase diagram for the carrier type in the thickness and Ca concentration space: n, n-p mixed, and p regime are depicted as red, pink and blue squares, respectively. Note that the phase space


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for the p regime shrinks as the film gets thin. (c) A cartoon showing the common trend of Rxy(B) depending on Ca concentration for all thicknesses. (i) Bi2Se3 is naturally n-type without Ca dopants. (ii) As Ca doping level increases, the (n-type) sheet carrier density gradually decreases, resulting in increased negative slope of the Rxy(B) curve. (iii-iv) With more Ca doping, the slope changes from negative (n-type) to positive (p-type) passing through a nonlinear n-p mixed regime.

Figure 3. Transport properties and morphology comparison in Ca-doped Bi2Se3 films. (a) 2D sheet carrier density (n2D) and (b) mobility (µ) vs Ca concentration of the 20 QL, Bi2Se3 films on BIS buffer layer with Se capping. The carrier densities are taken from the 20 QL Hall effect data of Figure 2 at B = 0 T. Data from severely nonlinear Rxy(B) curves (0.13% and 0.17% of Ca) are


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excluded due to their uncertain carrier densities. The carrier type changes from n to p, and finally to n, while the mobility decreases sharply right around the n-p crossing point and gradually as the Ca concentration increases further. (c, d) STM topography images of Ca-doped Bi2Se3 films at (c) x = 0.6% and (d) x = 6%. The morphology degrades significantly at the higher Ca concentration, where the carrier type of the film reenters into the n regime from the p.

Figure 4. Ca doping-dependent transport properties, around the charge neutral point, of 8 QL Bi2Se3 films capped by MoO3 and Se (at 300 mK). See Supporting Information Figure S4 and Section V for a broader doping dependence of a similar set of films. (a) Magnetic field dependence of Hall (Rxy) and sheet (Rxx) resistance with different Ca doping levels. (b) 2D sheet carrier density (n2D) and (c) mobility (µ) with different Ca doping levels. We used (dRxy(B)/dB)B = 0T

for 0% ~ 0.11%, and (dRxy(B)/dB)B = 15, 20T for 0.12% and 0.14% respectively, to extract net

carrier density and mobility from the nonlinear Hall curves while avoiding the n-p mixing and


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the quantum Hall effects. Note that the mobility sharply decreases when the carrier type changes from n to p. (d) Schematic layer structure of the 8 QL Ca-doped Bi2Se3 films grown on the BIS buffer layer and capped by MoO3 and Se. The pure Bi2Se3 data (0% of Ca, the leftmost data in ac) are taken from Ref.6.


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AUTHOR INFORMATION Corresponding Author *E-mail: [email protected] Present Addresses (N.K) Department of Physics, Massachusetts Institute of Technology, Cambridge, MA 02139, U.S.A. Author Contributions J.M., N.K., and S.O. conceived the experiments. J.M., N.K., M.S. grew the films, performed the transport measurements, and analyzed the data with S.O. W.Z. and W.W. performed STM measurements. The manuscript was written by J.M. and S.O. with contributions of all authors. Notes The authors declare no competing financial interest. ACKNOWLEDGMENT We thank N. Peter Armitage for comments and Pavel P. Shibayev for proofreading. This work is supported by Gordon and Betty Moore Foundation’s EPiQS Initiative (GBMF4418) and National Science Foundation (NSF) (EFMA-1542798). Transport measurements in National High Magnetic Field Laboratory were further supported by NSF (DMR-1157490) and the State University of Florida. ASSOCIATED CONTENT Supporting Information.


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Characterization of In2Se3-(Bi0.5In0.5)2Se3 Buffer Layer, control and estimation of Ca doping level in Bi2Se3 films, determination of carrier density and mobility in the n-p mixed regime, via two carrier model, comparison of alkaline earth metals as hole dopants: Mg, Ca, Sr, and Ba, and transport properties of another set of 8 QL Ca-doped Bi2Se3 thin films capped by MoO3 and Se, covering a broader doping range than Figure 4. This material is available free of charge via the Internet at http://pubs.acs.org.


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Solution to the Hole-Doping Problem and Tunable Quantum Hall

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